This Review describes behaviors and mechanisms governing heteroepitaxial nucleation and growth of group III (Al, Ga, and In)–selenium (Se) based semiconductors by molecular beam epitaxy and the properties of the resultant nanoscale films. With nine bonding electrons per pair, these chalcogenide semiconductors crystallize in a variety of locally tetrahedral bulk structures that incorporate intrinsic vacancies (atom-sized voids) lined with doubly occupied lone-pair orbitals, including layered, defected zinc blende and defected wurtzite structures. During heteroepitaxial growth, the choice of how the vacancies order and which phase results, as well as interface reactions, intermixing, surface passivation, and film morphology, are controlled by electron counting, substrate symmetry, and size mismatch. Nucleation and growth of , , and compounds on Si and GaAs, including initial reactions, layer nucleation, symmetry, crystal structure, defects, dimensionality, and stoichiometry, were studied with a combination of techniques, including photoelectron spectroscopy, x-ray photoelectron diffraction, scanning tunneling microscopy, x-ray absorption spectroscopy, and low energy electron diffraction. The unique crystal structure of was also investigated as a novel platform for doping with transition metals to create a dilute magnetic semiconductor: is ferromagnetic at room temperature, while results in the precipitation of MnSe. The present study provides new insight into growing interest in variable dimensional materials, using group III selenides as prototypes, to address the basic physical chemistry governing the heteroepitaxy of dissimilar materials.
Semiconductors based on compounds of group III and group VI elements are receiving increased attention worldwide. These elements are found as either constituents or dopants in more traditional semiconducting materials such as group IV (Se, Ge), III–V (GaAs, AlAs, etc.), or II–VI (ZnSe, CdTe, etc.) materials. The III–VI compounds are chemically compatible with more traditional semiconducting materials, and they exhibit similar optical bandgaps and lattice constants, as shown in Fig. 1.1,2 On the other hand, III–VI compounds exhibit many distinct structural, electronic, and optical properties that are not found in the traditional group IV, III–V, or II–VI materials. This combination of features holds significant promise for the development of semiconducting III–VI compounds as new device materials once issues related to heteroepitaxial growth on silicon and gallium arsenide substrates can be mastered. Layered III–VI materials also hold promise for isolated two-dimensional (2D) device applications.
The III–VI semiconductors comprise two principal stoichiometries: (a) compounds, such as GaS, GaSe, and InSe, and (b) compounds such as , , and . In addition, compounds such as exist as molecules and may be possible from the phase diagram; GaSe evaporates as .3 compounds exhibit a wide variety of crystal structures: the Crystallography Open Database2 lists over two dozen distinct structures and polytypes for and crystals where A = (Al, Ga, In) and B = (S, Se, Te). Selected crystal structures for and semiconductors are shown in Fig. 2.
Bonding states built from s and p valence states have room for eight electrons per two atoms; III–VI atom pairs have nine valence electrons. This imbalance in electron counting for III–VI materials leads to bulk crystals containing atomic-scale voids lined with doubly occupied lone-pair orbitals. Group III-Se semiconductors crystallize in a variety of structures that incorporate neutral intrinsic vacancies relative to a parent zinc blende or wurtzite structure, where the vacancies may form lines [Figs. 2(b)–2(d)], planes [Figs. 2(a) and 2(f)], or be distributed throughout the crystal [Fig. 2(e)]. This leads to the orbitals lining these voids forming 1D, 2D, and 3D electronic structures.
compounds typically present a layered structure, exhibiting weak van der Waals bonding between separate, covalently bonded B–A–A–B layers [A = Ga, In; B = S, Se, Te; Fig. 2(a)]. These materials are often referred to as “van der Waals materials” and the ability to isolate single B–A–A–B layers holds promise for the development of two-dimensional electronics.4 Layered bulk crystals exhibit strong optical and electrical anisotropy5–7 and high nonlinear optical coefficients in the infrared ranges, making them candidates for second harmonic generation materials.8–13 This interest led to a vital activity in the 1970s and 1980s to fabricate bulk single crystals of GaSe and InSe, and many optical and electrical properties were investigated.14
Compounds of stoichiometry are more common, since the valence for group III and VI elements are +3 and −2, respectively. Chalcogenide-based materials, however, are considerably different from oxides such as and . The lower ionicity of chalcogenides compared to oxygen leads them to prefer tetrahedral over octahedral bonding. They form a defected zinc blende or wurtzite structure, with one-third of the cation sites empty; for example, can be thought of as zinc blende ZnSe with every three atoms replaced by two atoms plus a vacancy. Different orderings of these empty sites along lines, helices, and planes lead to a variety of crystal structures (Fig. 2). These intrinsic vacancies, or atom-sized voids, allow to maintain semiconducting properties to a high level of impurity concentration15 and, in the case of Cr-doped , generate room temperature ferromagnetism.16 The crystal structure of these materials also explains their high radiation stability.17 Lithium atoms can fill these intrinsic voids in nanowires to generate new types of vacancy/lithium-atom ordered superlattices, demonstrating materials’ use in lithium ion storage, photovoltaics, and phase change memory.18 This is similar in concept to Li intercalation of a layered material, but Li resides on a helix instead of a plane. The intrinsic local electric field in layered makes it a potential photocatalyst for splitting water.19
Initial interest in materials, in particular , stemmed from the interfacial reaction inherent to II–VI on III–V heteroepitaxy, where interface disruption caused by formation during the initial stages of ZnSe/GaAs interface formation was found to be a severe obstacle.20,21 However, is a potentially useful material in its own right, especially in the thin film form. can passivate GaAs(001)22 and has strong potential for use in optoelectronic devices that exploit its visible bandgap.23 When the vacancies are ordered, has large absorption anisotropy .24 The close lattice matching of zinc blende to Si, GaP, and ZnS substrates (see Fig. 1) promotes its use in heterostructures. In particular, the close lattice matching (0.1%) and minimal interdiffusion (see below) at a properly formed /Si interface suggests its use in silicon-based electronics, where the bandgap in the epitaxial films25 lies in a convenient range for visible optoelectronics. The use of n-doped alloy deposited on p-Si heterojunctions for solar cells has been proposed.26
Heteroepitaxy of III–VI materials is essential for incorporation of these novel materials into complex device structures. These materials also serve as vital prototypes addressing the basic physical chemistry governing heteroepitaxy of dissimilar materials. This Review addresses the structure and thin film growth of -Se compound semiconductors using molecular beam epitaxy, primarily on silicon substrates, as well as the nature of bonding in bulk III–VI materials as it relates to controlling heteroepitaxy and the incorporation of transition metals for potential magnetic applications. The emphasis is on synthesizing work performed in collaboration between the Department of Physics and the Department of Materials Science and Engineering at the University of Washington. We both summarize work published in the archival literature and present work that has previously only been published in the dissertations of our graduate students.
II. CHEMICAL BONDING AND ITS INFLUENCE ON CRYSTAL AND ELECTRONIC STRUCTURE
A. AIIIBVI compounds
In layered compounds, each structurally identical layer is composed of single planes of group VI atoms on either side of a double plane of group III atoms [Fig. 2(a)]. The basic building block is formed by stacking four hexagonal monoatomic sheets in a B–A–A–B sequence with trigonal prismatic symmetry. Different stacking and/or rotation by 180° of these covalently bonded layers into a superlattice results in several polytypes, only some of which exhibit inversion symmetry. Polarization-dependent linear and nonlinear optical properties of layered , therefore, depend on the polytype.27
The B–A–A–B sequence of the quasi-two-dimensional layered semiconductors results in an unusual combination of chemical bonding types and hence electronic states.28 Interlayer bonding is via weak van der Waals interactions, with strong intralayer interactions. The cation–cation bond is symmetric (covalent), while the cation–anion bond has a partial ionic character, as is evident from the total valence charge densities calculated using empirical pseudopotentials by Depeursinge for GaS, GaSe, and InSe.29 In all three materials, valence electrons are strongly localized around the atoms, and the two atoms are linked by an -like bonding charge. Increasing the ionicity from GaSe (fractional ionicity of 0.66) to GaS (0.74) and InSe (0.80) increases the ionic character of the bond, pulling the electronic charges more toward the atoms. In addition, when Ga is replaced with In or Se is replaced with S, the atomic -states of the atom become less separated from the -states, leading to delocalization of bond charge and weakening of the central bond (increasing the bond length). Extending this further results in InS not forming a layered structure, but rather a three-dimensional orthorhombic structure, because the atomic In 5s-states are at higher energy than the S 3p-states, changing the charge transfer.
The electronic states within the B–A–A–B layer as measured with angle-resolved photoemission for GaSe are well described with a simple linear combination of atomic orbitals.30 Adjacent layers are weakly coupled, so the main difference between a single B–A–A–B quadlayer and the bulk crystal is a small splitting of the Se lone-pair states. Tight binding calculations show the bond is primarily of cation s and character, while the three bonds per A or B atom primarily involve states (four electrons per anion, two per cation); the remaining anion electrons reside in doubly occupied lone-pair orbitals on each surface of the layer. In GaSe, the hybridization of the Se orbitals is between that of and : in pure bonding, all bond angles would be 109°, while for pure , the intra- and bilayer angles would be 90° and 126°; in layered GaSe, the bond angle is about 100° and that between the central bond and any bond is about 118°.31 As discussed below, both the electronic30 and atomic32,33 structure of one half of the quadlayer stays essentially the same when the central bond is replaced by an bond during heteroepitaxial growth.
Semiconductors of the type typically crystallize in a structure based on zinc blende or wurtzite, in which one-third of the cation states are vacant, although layered or spinel structures are also reported. The structural vacancies order in different ways—in lines, helices, or planes—depending on both growth conditions and the choice of A or B atom. The use of the term “vacancy” is relative to the more familiar or zinc blende or wurtzite structures; it refers to neutral, atom-sized voids in the monoclinic unit cell. We also use the standard notation of the parent zinc blende and wurtzite crystals when referring to crystal surfaces and directions.
The bulk crystal structures derived from x-ray diffraction for zinc blende-based (Refs. 34 and 35) and wurtzite-based (Ref. 36) are shown in Figs. 2(b) and 2(d), respectively. The structures differ by the  or  stacking pattern of the arrays of vacancy lines, which are parallel to or . In , the defected wurtzite structure has vacancies aligned along a helix with an axis parallel to , as shown in Fig 2(e).37 In thin film growth on Si(001), vacancies in are aligned along lines [Fig. 2(c)].38 In these structures with linear or helical vacancy alignments, each cation (A) is bonded to four B atoms, while one-third of the anions have two cation neighbors and two-thirds have three. Under some conditions, forms B–A–B–A–B layers, with van der Waals bonding between them, as seen in , where the layers have hexagonal symmetry [Fig. 2(e)],39 or in epitaxial on GaAs(001), where the cubic layer stacking is in the  direction (not shown).40 The planar structure in Fig. 2(e) is similar to that of the topological insulator ;/ superlattices are of interest to explore quantum size effects in topological insulators.41 The intrinsic vacancies in can lead to passivated surfaces (Sec. III B), since termination does not result in partially occupied, “dangling-bond” surface states, as well as extensive local structural rearrangements around impurities without broken bonds (Sec. VI B).
III. ELECTRON COUNTING AND SURFACE PASSIVATION
Potential commercial application of chalcogenide III–VI semiconductors requires heteroepitaxial thin films, since the intrinsic vacancy structure of bulk crystals generally leads to poor mechanical and thermal properties. Growth of these polar materials on a nonpolar substrate (such as silicon) requires understanding of the electron counting at the interface and any modification or interdiffusion required to reduce interface dipoles42 and/or passivate the surface dangling bonds.43 Two important contributions related to electron counting are (a) the number of electrons available per atomic orbital and (b) the electrostatic potential difference generated when electrons transfer between atoms of different valence to fill those orbitals. Elimination of partially filled orbitals when bonds are broken at a semiconductor surface is a major driver of surface reconstruction; reduction of local electric fields at an interface between atoms of different valence is a major driver of interface mixing.
When III–VI materials, with nine electrons per atom pair, form a heterointerface with tetrahedrally bonded group IV, III–V, or II–VI materials, with eight electrons per pair, the imbalance in electron counting can lead to surface passivation (by fully occupying dangling-bond orbitals) and/or interface mixing. As described above, layered materials are of the form B–A–A–B, which may be thought of as two stable A–B bilayers that each have nine electrons per repeat unit [box in Fig. 3(d)]. One electron from each bilayer combines with one from the other bilayer to form the central A–A bond, six electrons from each bilayer form three A–B bonds, and the remaining two electrons fully occupy a “lone-pair” orbital on the layer surface [Fig. 3(d)]. This stable nine-electron bilayer with a single bond to another entity is ideal for passivating the surface of a tetrahedrally bonded semiconductor with eight electrons per unit cell when it is terminated in the plane perpendicular to a bond, i.e., the (111) surface of diamond or zinc blende or the (0001) surface of wurtzite.
A. Passivation of Si(111)
In bulk silicon, each atom is tetrahedrally bonded along the directions. The surface created by cutting perpendicular to a bond direction thus has a singly occupied “dangling bond” for each surface atom. To lower its energy, the Si(111) surface reconstructs to a complex structure with a unit cell involving a combination of dimer formation, adatom adsorption, and stacking faults that result in only 19 dangling bonds instead of 49 per unit cell.44 This complex structure is four layers thick and complicates heteroepitaxial growth on the Si(111) surface. It is possible, however, to remove the driving force for reconstruction by adding an electron to the dangling bond orbital. This may be accomplished, for example, by terminating the surface with hydrogen [Fig. 3(a)]. Simply exposing the Si surface to hydrogen does not remove the reconstruction, but preparing an Si(111) surface chemically in buffered HF leads to a nonreconstructed, H-terminated surface where each valence electron of a surface Si atom is in a covalent bond (three with Si, one with H).45 Note that the surface bilayer has nine valence electrons (four from each Si, one from H) per surface unit cell and has a single bond perpendicular to the surface to the bulk crystal below.
An alternate method for passivating the Si(111) substrate and removing the reconstruction is to replace the top layer Si atoms with As, which has one additional valence electron, by cooling an Si surface in an As flux from above 800 °C.46,47 The bilayer also has nine valence electrons per surface unit cell, four from Si and five from As [Fig. 3(b)]; three of the As valence electrons bond with the underlying Si, and the remaining two occupy a “lone-pair” orbital at the surface. The Si(111):As surface is extremely stable, with no permanent changes after atmospheric exposure;46 its formation during GaAs heteroepitaxy on Si(111) precludes laminar growth of GaAs, since its low surface energy relative to GaAs results in GaAs island formation.48
A thought experiment that moves a proton from Si to As at the Si(111):As surface creates an AlSe bilayer with a very similar electronic structure except for the dipole. GaSe termination of Si(111) has the same electron counting [Fig. 3(c)]. Exposure of Si(111) to a flux of either gallium selenide32,33 or aluminum selenide49 leads to such a bilayer termination of the Si(111) surface, despite the absence of a stable layered crystal structure for bulk AlSe. The stable, unreconstructed Si(111):GaSe structure is extremely interesting in its own right and is discussed in Sec. IV.
The atomic structure of bilayer-terminated Si(111) is very similar to half of a bulk GaSe quadlayer.30 Figure 4 shows the photoelectron diffraction results from Si(111):GaSe (Refs. 32, 33, and 50) and Si(111):AlSe.49 Photoelectron diffraction yields element-specific structure through scattering of photoemitted electrons by neighboring atoms and is well modeled by multiple-scattering calculations,51 enabling extraction of structural parameters. At high kinetic energies [Fig. 4(a)], the largest feature comes from forward scattering of Ga emission by the surface Se atoms; Se emission shows only weak first order diffraction peaks from in-plane Se–Se scattering.32,49 The dashed line in Fig. 4(a) highlights the 63° (65°) angle between the vertical and the Al–Se (Ga–Se) bond, which is close to the bulk GaSe value of 62°. At lower kinetic energies (KE),33 electrons experience significant backscattering; a stereographic projection of emission from Ga [Fig. 4(c)] shows peaks due to both the forward scattering from three neighboring Se at 65° and backscattering at normal emission from the underlying Si. The location of the interface Ga directly above the interface Si (top site), as opposed to three threefold hollows above the second (T4) or fourth (H3) layer Si, as well as the bond length, were determined by comparing multiple-scattering calculations to the interference pattern of Ga emission along the surface normal as the photon energy, and hence the electron wavevector, is varied [Fig. 4(e)]. Emission from Se atoms shows both backscattering from the three underlying Ga atoms at 65° and diffraction rings from both Ga and the six in-plane Se next-nearest-neighbors [Fig. 4(d)]. Varying the photon energy for Se emission along the Ga–Se bond [Fig. 4(f)] yields that bond length as well. These structural results as well as those from similar measurements on Si(111):AlSe (Ref. 49) are summarized in Fig. 4(b).
The similarity in atomic structure among Si(111):As, Si(111):GaSe, Si(111):AlSe, and GaSe (Fig. 3) continues to their electronic structure. Each exhibits a fully coordinated surface with a doubly occupied lone-pair state at or near the valence band maximum that disperses downward from the center of the Brillouin zone, as measured with angle-resolved photoemission spectroscopy.30 The states for a single bilayer of GaSe on Si(111) show a strong similarity to that of bulk GaSe, with the bond states having a similar character to the states, and the Se lone-pair orbitals roughly the average of the bulk GaSe states that are split due to layer–layer interactions. The near-surface bands in Si(111):As, on the other hand, more closely resemble those of bulk Si.47 The GaSe bilayer alters the Si bands for several atomic layers, reflecting strain, surface dipoles, or some other long-range effect. However, core-level photoemission shows the interface Si bonded to Ga to be in a bulklike environment. No interface shift (<0.1 eV) of Si was resolved, unlike the 0.75 eV shift of the Si(111):As system. This serves as a reminder of the complex combination of initial and final state effects controlling core-level energies near a surface, as well as how surface band dispersion and charge transfer are separate issues. The similarity in the lone-pair dispersions between the bilayer and bulk crystal highlight the two-dimensional nature of the electronic states in bulk layered GaSe.
Despite the strong similarity in atomic and electronic structures of Si(111):AlSe and Si(111):GaSe, their resistance to contamination is quite different. As shown in Fig. 5, GaSe-terminated Si is much less reactive than Si(111):AlSe.52 When exposed to atmospheric-pressure air, nitrogen, oxygen, or nitrogen gas saturated with water vapor, Si(111):GaSe shows no change in surface-sensitive core-level photoemission [see left side of Figs. 5(b)–5(e) for result with atmospheric exposure]. Valence band emission [Fig. 5(a)] is suppressed by physisorbed species after atmospheric exposure (dashed line), but a 1 min vacuum anneal at 480 °C restores the system to the initial conditions [dotted line (green online)], including the Se lone-pair surface state near (arrow). On the other hand, Si(111):AlSe permanently reacts with , water vapor, or atmosphere at room temperature, exchanging O for Se. This is evident in the large (70%) decrease in Se intensity [Fig. 5(c)], large O emission [Fig. 5(e)], and the addition of a high binding energy component on Al due to the higher electronegativity of oxygen relative to selenium [Fig. 5(b)]. The valence band structure permanently changes to open a wide bandgap upon exposure [Fig. 5(a)]. The oxidized AlSe bilayer does, however, protect the underlying silicon from reacting with the atmosphere as well as a GaSe bilayer does [Fig. 5(d)].
In contrast to AlSe and GaSe, InSe does not form a stable bilayer on Si(111),53 despite the existence of a bulk layered InSe structure similar to that in Fig. 2(a). While exposure of Si(111) to evaporated GaSe or results in an initial 1:1 ratio of cation to Se, even if additional Se or Al flux is supplied, respectively, InSe evaporation results in approximately one-third monolayer of Se deposited before In starts to stick. Subsequent growth is of , as discussed in Sec. V C.
B. Passivation of Si(001)
Creating a (001) surface of silicon breaks two bonds per surface atom, and the native surface reconstructs to form dimers in the or direction (alternating with each layer), reducing the number of dangling bonds by half.54 The Si reconstructed surface is highly reactive, and when exposed to an Se flux [as in growth of ZnSe (Ref. 55) or (Ref. 50)], it forms SiSex similar to the Si formed when it is exposed to oxygen.56 This reaction can be prevented by terminating the surface with arsenic. With five valence electrons, As forms dimers similar to the Si(001) surface, but instead of dangling bonds, each surface As exhibits a doubly occupied lone-pair orbital.57 This passivated Si(001):As surface, which leads to islanded growth of GaAs on Si(001),48 enables laminar heteroepitaxial growth of group III selenides, as discussed in Sec. VI A. Exchanging the outermost Si–As layer to create a stable Ga–Se bilayer on Si(001) in a manner similar to that discussed above for Si(111) has not been observed. A few bilayers of deposited on Si(001):As, however, form a stable structure resistant to oxidation.
The formation of at the surface of Si(001) is the foundation of microelectronics. However, as devices shrink, a higher dielectric constant insulator is required to maintain a reasonable capacitance. Titanium dioxide, which in its anatase crystal structure is nearly lattice matched to silicon, is a candidate high dielectric constant material. However, its exploitation is hindered by the formation of a low-dielectric-constant interface layer that dominates the system capacitance.58 This problem can be remedied through the use of surface passivation, namely, Si(001):As:.59 In contrast to clean Si, Si(001):As is stable when exposed to , even at 450 °C; the black line in Fig. 6 shows a silicon core level (b) and valence band emission (d) identical to clean Si(001):As. When reactive Ti is added to the oxygen, however, substrate oxidation occurs for Si(001):As, with reacted components present in both the Si [Fig. 6(b), dots (blue online)] and As (not shown) emission. On the other hand, a 0.6 nm buffer layer of on top of Si(001):As prevents oxidative disruption of the interface and enables laminar growth of without the formation of [note the absence of oxide peak for Si in Fig. 6(a)].59 The valence band offset between Si and is similar both with [Fig. 6(c)] and without [Fig. 6(d)] the buffer layer, with the Fermi level pinned near the top of the oxide bandgap [solid lines (red online)] and the bottom of the silicon gap (dotted lines). The band alignment was deduced from knowing the bulk separation between the valence band maxima and the levels for Ti and Si, as well as the bulk bandgaps.59 Unfortunately, this staggered type-II band alignment is not appropriate for using anatase as a gate insulator for n-type silicon. The structure, growth, and properties of this buffer layer are detailed in Sec. VI.
C. Passivation GaAs(111) and GaAs)
The equivalent surface of GaAs to Si(111) can be either Ga-terminated or As-terminated . These polar surfaces have an intrinsic dipole as well as dangling bonds, leading to a complex reconstruction that is different on the two surfaces. On GaAs , removing 1/4 of the surface Ga leads to a reconstruction with no partially filled orbitals; GaAs takes either a or reconstruction, depending on the As concentration.60
Electron counting arguments for GaAs surfaces are more complicated than those for silicon since the GaAs bond is not symmetric: in the bulk, each Ga–As bond averages of an electron from Ga and of an electron from As. At the ideal surface, there is thus electron per As dangling-bond orbital; to fill these orbitals, three electrons thus need to be added for every four surface anions. This may be accomplished by replacing three out of four atoms with . Combined photoemission and RHEED of GaSe deposition on GaAs is consistent with this picture, with a stable bilayer bound to the substrate, followed by subsequent deposition of stoichiometric, layered GaSe (see Fig. 7, right).61 At the ideal (111) surface of GaAs, the Ga-terminated surface will have its dangling bond orbitals each occupied by electron. Removing one atom out of four thus leads to a surface where each Ga–As bond has two electrons and the Ga-derived dangling bond orbitals are all empty. Photoemission spectroscopy61 shows that the deposition of GaSe on GaAs(111) results in extensive interdiffusion of As and Se, and ∼1 ML of Se is observed as an extra layer between the GaAs substrate and subsequent deposition of layered GaSe (Fig. 7, left). Figure 7(c) compares the GaSe/GaAs interface for the two surface terminations; note the different location and bonding of the Se interface layer as well as the enhanced interdiffusion on the Ga-terminated (111) surface. In both cases, the Se-passivation leads to a discommensurate interface between the GaSe and lattice-mismatched underlying substrate.61
IV. GaSe-terminated Si(111)
GaSe-terminated Si(111) is an extremely stable, unreconstructed surface that leaves the underlying Si in an essentially bulklike environment. This quasi-two-dimensional material is important in its own right as a prototype system to study the formation of domain boundaries (DBs), charge redistribution near substitutional defects, and spin–orbit effects in photoelectron diffraction.
As discussed in Sec. III, exposure of the reconstructed Si surface to evaporated GaSe completely removes the complex reconstruction. Figure 8 contains scanning tunneling microscopy (STM) images of Si(111) with [Figs. 8(b) and 8(d)] and without [Figs. 8(a) and 8(c))] a terminating GaSe bilayer.62,63 On the nanometer scale [Figs. 8(c) and 8(d)], the addition of GaSe reduces the surface corrugation from hundreds to tens of picometers as adatoms and deep corner holes are eliminated ( unit cell denoted by the gray (green online) rhombus is 2.7 nm on a side). The Si(111):GaSe images are very similar to STM of Si(111):As.64 Atomic corrugations in GaSe bilayer-terminated Si(111) can be visualized with both positive (empty state) and negative (occupied state) bias STM, with the peaks in tunnel current at the same location (above Se atoms) in each case.62 The Si(111):GaSe island, pit, and corrugated step-edge structure apparent in Fig. 8(b) results from the displacement of Si atoms when the structure, which is four layers deep and has four more atoms per unit cell than the ideally terminated surface, reacts to form the bilayer-terminated surface.
Three types of intrinsic defects with different dimensionalities are observed with in situ STM for GaSe bilayer-terminated Si(111): zero-dimensional clustered point defects (CPDs), one-dimensional DB, and two-dimensional Ga-terminated Si regions (Ga/Si).62 DBs are formed during the coalescence of two GaSe bilayer domains whose orientations differ by 180°: one with Ga–Se bonds in same direction as Si–Si bonds (zinc blende or diamond stacking), and hence Se in the H3 (hollow) site, and the other with Ga–Se bonds antiparallel to Si–Si bonds (wurtzite or layered GaSe stacking), and hence Se in the T4 site directly above the second-layer Si. The rotated (wurtzite-orientation) domains are kinetically stable only when the films are grown at temperatures below 550 °C (which is the growth temperature for the single-domain GaSe films in Fig. 4); the local orientation is likely related to whether nucleation occurs on the faulted (GaSe stacking) or unfaulted (Si-stacking) half of the unit cell. These orientational domains are distinguished by an apparent height difference of ∼0.5 Å in the empty state imaging by STM [Fig. 9(a)] but have similar apparent heights in occupied states [Fig. 9(b)]. A model deduced from high resolution imaging of the empty [Fig. 9(c)] and occupied [Fig. 9(d)] states at a DB includes a Ga dimer structure similar to the Si dimers found at the boundary between faulted and unfaulted regions of the Si(111) unit cell, as well as occasional pairwise substitution of Se atoms by Si at the domain boundary that likely reduces the electrostatic energy of the Ga–Se dipole interaction across the boundary.62
Orientational domain boundaries frequently terminate in Ga-terminated Si regions, where Ga segregates under Se deficient conditions. The dark region just below the inset in Fig. 9(b) is an example. A larger Ga-terminated region is imaged in Fig. 9(e). The structure is similar to the structure reported for about 1 ML of Ga on Si(111).65 These domains, whose existence may also be detected through a low-binding energy component in Ga photoemission, are suppressed by including an extra source of Se during growth.50
Isolated point defects are not observed with STM for the GaSe bilayer; rather, defects appear as CPDs involving several unit cells. For example, the inset in Fig. 9(a) shows three 0.4 Å protrusions at Ga positions separated by three unit cells; these appear as 0.1 Å dents in occupied state imaging [inset in Fig. 9(b)]. Ohta et al.62 reported several distinct CPDs, all of which extend over at least three unit cells and propose that these CPDs are best explained by a localized charge density wave that scatters from a substitutional point defect and are evidence of intermixing or interdiffusion during growth. The CPD structures observed with STM differ for the two domain orientations. The CPD in the insets in Figs. 9(a) and 9(b) is the most common defect for the rotated domains; it is centered on the location of a Ga atom and is attributed to substitutional . The most common defect on nonrotated domains appears as a protrusion centered on an Se site with an additional depleted hexagon (sides of length two unit cells) in occupied state images that is not present in empty state images. Several may be seen in the upper-right of Figs. 9(a) and 9(b). In both of these cases, the STM image reflects nanometer-scale charge redistribution within the 2D bilayer in the presence of a 0D substitutional point defect.
The highly ordered structure of Si(111):GaSe, with its shallow Ga core level (binding energy ), provides a fruitful testing ground to investigate the role of spin–orbit interactions in photoelectron diffraction. In the case of As-terminated silicon, differences in the angular variation of intensity for As and emission were attributed solely to the difference in energy of the two states,66,67 while for In-terminated Si(100), additional structure in the In spin–orbit branching ratio (SOBR, ratio of to emission intensities) was attributed to unknown matrix element effects.68 Si(111):GaSe displays a surprisingly strong (>15%) angular variation of the Ga 3d photoemission SOBR.50,69 Figure 10(a) shows the angular variation along the azimuth for the full Ga emission [equivalent to horizontal diameter in Fig. 4(c)] and separately for the spin–orbit components, and . The sample was rotated while the angle between the linearly polarized incident photons and detected electrons was held at 60° in the plane of the photon polarization. The inset in Fig. 10(a) shows the distinct variation in the Ga peak shape as the spin–orbit ratio varies with the emission angle; the angular variation of the SOBR is shown in Fig. 10(b).
The origin of the variation in SOBR with the angle can be addressed through calculation by adapting relativistic photoabsorption matrix elements, originally developed for modeling x-ray absorption fine structures,70 and applying them to multiple-scattering analysis of photoelectron diffraction.50,69 Figure 10(c) shows results of such calculations using FEFF,70 which features an ab initio treatment of atomic properties based on a spinor-relativistic single-configuration Dirac–Fock atom code. The SOBR for Ga emission from Si(111):GaSe is shown both with (solid line) and without (dotted line) taking into account the 0.4 eV difference in their kinetic energies, assuming the structure of Fig. 4(b), as well as that ignoring scattering from neighboring atoms (dashed line). The results show that SOBR angular variation is not strongly dependent on the energy difference between the two spin–orbit components nor is it an atomic feature. Rather, the SOBR angular variation is the result of different diffraction conditions for the two spin–orbit components due to different final state wave functions; qualitatively, it arises from their different andwave decomposition. The calculations predict both the position and relative amplitude of the Ga SOBR variations. The absolute amplitude of the experimental effect is roughly four times larger than the theoretical prediction for Ga emission [compare Figs. 10(b) and 10(c)]. On the other hand, the variation of the SOBR (not shown) is about for both experiment and theory. The experimental enhancement of the Ga SOBR variation may be attributed to overlap between Ga and Se levels; these levels were not included in the theoretical calculation but would increase the overlap between the initial and final states for Ga. Another factor impacting both Ga and Se modeling is the use of spherical muffin-tin potentials for the calculation despite the presence of a crystal field, which may account for the higher average branching ratio in the calculation than in the data for both Ga and Se.50
V. GROUP III SELENIDE NUCLEATION AND GROWTH ON Si(111)
We turn now to the initial stages of heteroepitaxial growth of semiconductors on silicon. The wide variety of stable structures for III–VI materials means that the choice of substrate termination and orientation can determine the structure and stoichiometry of the III-Se film, and hence its potential uses. All the structures in Fig. 2 involve different stackings of hexagonal layers with threefold symmetry, which is the structure of the Si(111) surface. The (001) plane of silicon, on the other hand, has square atomic arrangements with two perpendicular mirror planes. Of the common crystal structures, only (defected zinc blende) is cubic, although there is71 a high pressure defected-spinel phase for ; the other structures are mainly hexagonal. This raises the question of whether heteroepitaxial growth of group III selenides on Si(001) and Si(111) will lead to different structures or if weak van der Waals bonding can lead to discommensurate growth of hexagonal layered materials in either case. Also, the bulk cubic phase exhibits vacancy lines along the direction, which is perpendicular to , but not to ; this leads to different growth dynamics and vacancy ordering on the two surfaces. In this section, we discuss the nucleation and growth of aluminum, gallium, and indium selenide on Si(111); Sec. VI discusses aluminum and gallium selenide growth on Si(100).
Aluminum selenide first forms the Si(111):AlSe bilayer termination discussed in Sec. III, despite not having a bulk layered form similar to Fig. 2(a); it then continues growth along as defected wurtzite , with .72 Gallium selenide, on the other hand, continues growth on the Si(111):GaSe bilayer termination as layered GaSe.73 Indium selenide, which does exhibit a bulk InSe layered structure,74 does not form bilayer termination; it first terminates the Si(111) surface with a partial monolayer of Se and then continues as , although with a complex surface reconstruction.53,75 The initial film morphology for deposition of on Si(111), as imaged with scanning tunneling microscopy, is shown in Fig. 11 for evaporation source materials of (top), GaSe (middle), and (bottom).
A. growth on Si(111)
Exposure of Si(111) to evaporated GaSe at substrate temperatures in the range 400–585 °C results in essentially stoichiometric GaSe deposition throughout, except for a small fraction of, the Ga-terminated regions [e.g., Fig. 9(e)] that may be eliminated by adding extra to the flux to counteract its smaller initial sticking coefficient . Initially [Fig. 11(d)], GaSe forms small islands that later merge to form a flat bilayer, followed by subsequent nucleation and growth of layered GaSe once the bilayer is complete.63 Photoemission reveals the ratio of Ga to Se to remain constant (and equal to that of bulk GaSe) throughout this evolution. The sub-bilayer growth imaged in Fig. 11(d) shows small islands of height close to that of the smooth, completed bilayer visible along the left of the image; disordered structure is present between the islands. Figure 11(e) shows a film just after completion of the GaSe bilayer; the bright islands are 0.8 nm high, which is the height of the Se–Ga–Ga–Se quadlayer, easily distinguished from the 0.3 nm step height between bilayer-terminated regions [rendered as black and light gray (orange online)] that is characteristic of the silicon substrate.
Once the bilayer is complete, the GaSe growth rate decreases, with only about 10% of the atoms sticking above 450 °C; there are very few nucleation sites for the next layer. The main nucleation sites are at domain boundaries, with additional sites at substrate step-edges, as seen in the larger scale image in Fig. 12(a) and expanded view in Figs. 12(b) and 12(c). Figure 12(b) images a collection of triangular GaSe islands a few nm in size; this film was deposited at so that the bilayer exhibits many domain boundaries, and shows that most islands inhabit one of two orientations that differ by 180°. Figure 12(c) is the horizontal gradient of the image in Fig. 12(b), highlighting orientational domain boundaries; most islands show a domain boundary terminating under them (circles). With further deposition, growth continues as layered GaSe, with its smooth surface and 0.8 nm steps. Initially, the layers confine themselves to a single substrate terrace [Fig. 12(d)], but once the second layer nucleates, the flexible layers eventually “carpet” the substrate steps [inset in Fig. 12(e)].63
B. growth on Si(111)
The structure of sub-bilayer aluminum selenide growth on Si(111) [Fig. 11(b)] shows similar structures to gallium selenide deposition [Fig. 11(d)], although AlSe tends to form lines along the direction rather than the compact GaSe islands on the structure. In Fig. 11(b), the smooth triangle is a complete AlSe bilayer. Once the bilayer is complete, additional nucleation and growth is as defected wurtzite . In Fig. 11(c), the step height of the islands atop the AlSe bilayer is 0.3 nm, equal to a single hexagonal bilayer (HBL) of the wurtzite structure. The islands exhibit a structure consistent with 1/3 of the Al sites being vacant, this reconstruction continues as the surface as the film grows thicker.72
Surface-sensitive core-level spectroscopy, which probes interface chemistry and stoichiometry, is shown in Fig. 13 for (a) aluminum selenide72 and (b) indium selenide53 growth. Films were deposited while slowly moving a shutter across the substrate to create a wedge-shaped sample, and photoemission spectra were acquired as a function of position, and hence film thickness. The photoemission peaks for aluminum selenide growth [Fig. 13(a)] show that the bilayer component of Al and Se emission (shaded, dashed line) is steadily buried at the same rate as emission from the underlying silicon, indicating that the bilayer remains intact. The single Al peak and two Se components separated by about 1 eV that arise with additional HBL have the same energies as those for an 8 nm-thick film,76 showing that the bulk structure depicted in Fig. 2(d) starts as soon as the first bilayer is completed. High energy photoelectron diffraction of these emission peaks exhibits the symmetry and expected positions for wurtzite.76
C. growth on Si(111)
While AlSe and GaSe form a bilayer termination of Si(111), InSe does not.53 With a large lattice mismatch (see Fig. 1), indium does not readily fit into the Si structure. Instead of forming In–Si bonds, the growth initiates instead with Se replacing the adatoms in the Si(111) structure. Core-level spectroscopy [Fig. 13(b)] shows that a partial Se layer sticks to the surface before any indium, but the Se:In ratio quickly converges to 3:2 by the time one monolayer of In has been deposited. High resolution scanning tunneling microscopy for the Se-only layer [Fig. 11(f)] shows protrusions at the position of the Si adatoms; these protrusions appear taller in occupied state images than in empty state images, consistent with .
Further deposition of indium selenide on Si(111) in the temperature range = 570 –580 °C results in growth of defected wurtzite .53,75 Lower temperatures tend to result in disordered, Se-rich, rough films, while higher temperatures are In-rich with very little Se.77 The incomplete HBL coverage imaged in Fig. 11(g) shows atoms scattered seemingly randomly and occasionally clustering to form ordered pinwheel shapes with a layer height of 1.4 Å above the base level; the uncovered regions (dark triangle) exhibit a local reconstruction. This pattern is consistent with 1/3 ML of Se in adatom sites, while the atoms above these are likely indium. The 1.4 Å step height is consistent with adding In and then incorporating the underlying Se into a bilayer structure. Additional deposition leads to coalescence of the pinwheel shapes into hexagonal figure-eight shapes [Fig. 11(h)].
As more is added,53,77 the layers maintain the same surface structure with steps of height 0.3 nm, as expected for defected wurtzite , and not the 1.0 nm steps of layered or the 0.83 nm spacing of layered . Despite the large in-plane mismatch, does not form islands to relieve interfacial stress; rather, the hexagonal lattice constant of the layer is very close to its bulk value starting at the first layer, resulting in a discommensurate interface. The surface forms a canted honeycomb pattern oriented at angles of about ±2° with the substrate directions [Fig. 14(a)]. The Fourier transform of the STM image in Fig. 14(a) is shown in Fig. 14(b), forming a hexagonal pattern where each vertex is a pair of spots. Low energy electron diffraction (LEED) from this surface at 16.4 eV [Fig. 14(c)] shows the underlying hexagonal pattern plus superstructure spots consisting of two triangles; the triangles indicate a total of six domains that are nearly rhombohedral, but with lattice vectors of two slightly different lengths. Electron diffraction at 34.2 eV shows a more complicated pattern [Fig. 14(d)] with multiple superstructure components. The LEED data are well modeled78 with six domains (three left-handed and three right-handed) at ±2° from the three symmetry-equivalent directions on the surface, with lattice constants 1.72 and 1.84 times the underlying lattice vector of 4.14 Å ( and ).77
STM of several-nm-thick films of deposited at 575 °C shows that the complex, discommensurate reconstruction persists at the surface, while x-ray diffraction shows that the underlying film exhibits the defected wurtzite structure.53 When amorphous is deposited on Si(111) at room temperature and then annealed to 380 °C, it crystallizes to without the complex reconstruction; STM shows smooth terraces with step-height 0.3 nm and both LEED and STM exhibit a surface reconstruction from ordered In vacancies. The crystallized material can be returned to the amorphous state with a high energy laser pulse, indicating potential application of films for phase change memory.75
In summary, the evolution of nanostructure morphology and local chemical environment during heteroepitaxial growth of the group III (Al, Ga, and In)–selenides on Si(111) lead to a variety of structures and morphologies. Despite the strong similarity between AlSe and GaSe in atomic and electronic structure during the deposition of their first bilayers, subsequent growth is quite different. Once the bilayer is completed, nucleation and growth of AlxSey results in an alternating Al–Se–Al–Se stacking sequence consistent with the defected wurtzite structure of Al2Se3, whereas gallium and selenium exhibit the layered GaSe structure, with van der Waals bonding between covalently bonded Se–Ga–Ga–Se layers. In the case of indium selenide, an InSe bilayer is not formed, but it first terminates the Si(111) surface with a partial monolayer of Se, and then continues as , although with a complex surface reconstruction. Potential use of bilayer-terminated Si(111) for heteroepitaxy of other materials has not been fully explored, although its low surface energy makes it likely that materials which are not intrinsically layered will exhibit islanded rather than laminar growth, similar to growth48 of GaAs on Si(111):As. Attempts to grow GaSe on As-terminated Si(111) resulted in about two bilayers of a cubic interface layer, but the growth then switched to layered GaSe.63
VI. GROUP III SELENIDE NUCLEATION AND GROWTH ON Si(001)
A. Pure group III selenides
The growth of group III selenides on Si(001) is complicated by two factors: the square symmetry and the high reactivity of the (001) surface with selenium. The former leads to novel crystal structures constrained by the substrate symmetry, while the latter can be overcome through the use of arsenic passivation. As discussed in Sec. III, the same tendency of Si(001) to react with oxygen to form amorphous leads to a similar reaction with selenium, but passivation of the Si(001) surface via As-termination enables the growth of selenides without forming amorphous .
Scanning tunneling microscopy and surface-sensitive core-level spectroscopy for aluminum53 and gallium63 selenide growth on Si(001):As are shown in Figs. 15 and 16, respectively. The Si(001):As surface consists of rows of As dimers along directions that rotate by 90° with each layer [Fig. 15(a)]. The first selenide layer nucleates differently for and deposition: nucleates in lines that are perpendicular to the As dimer rows immediately beneath it [Fig. 15(g)], while nucleates parallel to the underlying dimer rows [Fig. 15(b)]. Photoemission spectroscopy of the underlying substrate shows that deposition does not dramatically change the line shape of the Si or As emission, indicating the absence of strong bonding or interdiffusion [Figs. 16(c) and 16(d)], while for deposition, the interface silicon changes to a more bulklike environment (suppression of the high binding energy peak associated with Si bonded to two arsenic atoms) and the As broadens to contain at least two components [Fig. 16(b)]. The narrow Si interface emission is indicative of forming Si–Ga bonds at the interface, which, similar to the (111) case, exhibit no chemical shift from the bulk, instead of Si–As bonds, which exhibit a 0.5 eV chemical shift. The As spectrum [Fig. 16(c)], as well as its photoelectron diffraction pattern,63 indicate substitutional resulting from diffusion into the first couple of cubic bilayers in the overlayer.
As growth proceeds, the morphology is most ordered after the deposition of one to two cubic bilayers (CBL). Narrow rods, several nm long, alternate direction with each substrate terrace step [Fig. 15(c) and 15(d)]; Si(001):As with dimer rows parallel to the nanorod structure is still visible between the nanorods, which accounts for the persistent photoemission intensity from the substrate [Fig. 16(d)] and As interface layer [Fig. 16(c)]. Se emission [Fig. 16(b)] from the first one to two CBL exhibits a peak of similar half-width to that seen for (111) growth of .72 Once the film is more than about two BL thick, however, the Se core-level loses structure and broadens, indicating a disordered film, and STM shows the nanorod structure to break up into smaller and smaller pieces that show much less alignment with the underlying substrate [Fig. 15(e)].53
In contrast to , the quality of heteroepitaxy on Si(001):As improves with increasing thickness. The first CBL coalesces into smooth, broad anisotropic islands with rough edges; the substrate dimer rows do not persist [Fig. 15(h)]. By about two CBL, elongated nanorods develop smooth edges and a common width of about 3 nm, while photoemission (Fig. 16, right) shows development of an interface dipole between the substrate Si [Fig. 16(d)] and the overlayer (overlayer peaks shift to higher binding energy faster than the Si); the overlayer contains single-component gallium [Fig. 16(a)] and dual-component selenium [Fig. 16(b)]. The As [Fig. 16(c)] integrated intensity decays with thickness more slowly than the Si [Fig. 16(d)], while photoelectron diffraction shows the As to have the same diffraction pattern as Se (and different from Ga),63 indicating As is substituting for Se in the structure and not all remaining at the interface. Once the film is three CBL thick, the nanorods sharpen to less than a nanometer wide. As growth continues beyond three CBL, the rods maintain this minimum spacing of three dimer rows (1.2 nm), and minimum height differences of 0.26 nm, the CBL spacing.38,63
High resolution imaging of the stable nanorod surface [Fig. 15(k)] shows 10 pm atomic corrugation with periodicity 3.9 Å, the bulk (second-neighbor) distance between adjacent Ga or Se atoms.38 Comparison of empty and full state images, as well as photoelectron diffraction and comparison of photoemission intensities at different photon energies (and hence escape depths), shows the surface to be Ga terminated.63 These one-dimensional nanorods lead to the formation of one-dimensional vacancy rows in the film as the growth proceeds [Fig. 15(f)]. Angle-resolved photoemission shows a one-dimensional electronic state along  near the valence band maximum that is evidence for vacancy alignment in this direction throughout the film;79 this state is similar to calculations80 of Se lone-pair states lining the vacancy rows in .
The striking difference between growth morphologies of and on Si(001) is strongly connected to the behavior of arsenic at the interface.53,63 Given that the nanorods are cation terminated, the initial alignment of the nanorods parallel to the underlying As dimers requires that aluminum atoms are inserted into the As dimer bond for a zinc blende stacking arrangement [Fig. 17(a)]. At the level of one or two CBL, the resulting growth, with Al vacancies aligning in directions, is stable. However, with further deposition, the field arising from the large dipole associated with the intact polar/nonpolar interface leads to instability, which, when combined with the tendency for to take a wurtzite rather than zinc blende structure, results in the breakup of the nanorod structure. For growth, on the other hand, the exchange of As with Ga at the interface [Fig. 17(b)] leads to a nonpolar bond (Si shows no chemical shift), and the electron counting resulting from a combination of Ga vacancies and arsenic interdiffusion allows the surface to create and maintain a structure with no net dipole and no dangling bonds. For the first two to three layers, As interdiffusion leads to fewer vacancies (have the same electron counting as one ), which suppresses the nanoridge formation; once the As supply is depleted, the aligned nanoridge structure stabilizes. The stability of this structure and the absence of partially filled surface orbitals account for the extreme resistance of this structure to oxidation, even in the presence of titanium at elevated temperature, as discussed in Sec. III B.
B. TM–doped Ga2Se3 on Si(001):As
Commercial application of semiconductors relies on doping with impurities. The intrinsic vacancies inherent in semiconductors raise the question of how dopants may incorporate in these crystals. For example, a group-II dopant atom could replace a group III atom, and thus act as an acceptor, or occupy a vacancy, and thus act as a double donor. Besides electronic considerations, intrinsic-vacancy materials raise interesting scientific questions related to the potential mechanisms for ferromagnetism in dilute magnetic semiconductors, namely, mediation by free carriers [as in Mn:GaAs (Refs. 81 and 82)] or stationary defects (as in ).83 In GaAs, Mn doping at the level of a few percent leads to low-temperature ferromagnetism, with Mn contributing both the localized spins (five unpaired electrons) and the itinerant holes that couple them [two bonding electrons where Ga would have three ()].81 In Co-doped anatase , ferromagnetism has been observed in insulating films, with the coupling between Co spins attributed to localized defect states.83 With defect-associated states at the top of the valence band, namely, the Se lone-pairs that line vacancy rows,79,80 has potential for both mechanisms to contribute. Density-functional calculations of manganese, chromium, or vanadium inserted into a vacancy site in suggest that any of these dopants may result in spin-polarized states, with the largest net spin near the Fermi level predicted for Cr.84,85 These calculations, which were for a structure that nature did not choose to make in our experiments, nevertheless inspired exploration of Mn and Cr doping at the level of a few atomic percent in epitaxial films for potential magnetic applications. As discussed below, epitaxial is not a promising dilute magnetic semiconductor, but epitaxial exhibits room temperature ferromagnetism. Both dopants, however, have significant impact on the heteroepitaxial growth of .
Mn doping of films on Si(001) leads to disproportionation into lattice-matched MnSe islands and a low-Mn-concentration laminar film before the doped film is either thick enough or at a high enough concentration of Mn to have significant magnetic response.86,87 Chromium doping, on the other hand, leads to room temperature ferromagnetism.16 The unique vacancy structure of allows Cr to incorporate in a locally octahedral environment without major distortion of the overall lattice, while the Cr states hybridize with the Se lone-pairs, resulting in laminar, ferromagnetic films.
1. Mn-doped Ga2Se3
Manganese is soluble in bulk up to in , after which there is first precipitation of the quasi-zinc blende compound , and then, for disproportionation into and rock salt MnSe, although zinc blende-based is observed to higher Mn concentrations above 930 K.88 For growth on Si(001), however, both zinc blende and rock salt MnSe are better lattice-matched to the substrate than either or the bulk substitutional alloy with . This leads to disproportionation into MnSe islands during attempts to fabricate Mn-doped films on Si(001).86,87
The morphology of Mn-doped films deposited on Si(001):As varies with both thickness and Mn concentration, and to a lesser extent on the presence or absence of an undoped buffer layer (see Fig. 18).86,87 Mn concentration in Fig. 18 was determined from the incident flux, measured as , where QCM denotes the room temperature quartz crystal monitoring the flux. At low Mn concentrations and/or thickness (e.g., sample D1 in Fig. 18, which was deposited on a 4° miscut wafer to align domains), films are laminar with a nanoridge structure similar to that described above for pure films [e.g., Fig. 15(j)]. As the amount of deposition and/or Mn concentration is increased, there is a transition to an islanded morphology, with flat rectangular islands protruding from terraces that exhibit a distribution of nanoridge spacings and orientations similar to pure films; the specific island morphology varies with the growth parameters (compare D7 and D8 in Fig. 18), but island edges are generally aligned with and directions. An intermediate “platelet” regime (e.g., sample D5 in Fig. 18) has tall narrow islands nucleated on terraces that contain platelets of about 30 oriented and uniformly spaced nanoridges. The platelet morphology is metastable, converting to the islanded structure upon a 30 min anneal at the growth temperature of 500 °C. Samples with an buffer layer of pure exhibit islands with a somewhat larger footprint (compare D8 and B4 in Fig. 18), but the local atomic structure of both the islands and the terraces between them are similar with and without a buffer layer.
Information deduced from STM, extended x-ray absorption fine structure (EXAFS), x-ray absorption near-edge structure (XANES), and photoemission indicates that the tall, rectangular prism islands are rock salt MnSe.86,87 In situ STM shows the islands to be atomically flat and unreconstructed, exhibiting occasional 0.27 nm height steps characteristic of MnSe; STM also shows empty-state “protrusions” associated with Mn electronic states in the nanoridges on terraces between islands. Mn L-edge XANES and photoemission indicate that Mn is in the oxidation state both before and after air exposure, with more surface oxidation of Mn in laminar films than in islanded morphologies. Mn, Ga, and Se K-edge EXAFS are best fit by rock salt MnSe islands (when present) plus Mn substitution into in flat terraces, whether as a laminar film or the region between islands (some of which oxidizes in air-exposed samples).
Rock salt MnSe has an fcc Se sublattice that is essentially identical (0.2% mismatch) to that in , and the island orientation observed with STM is consistent with a continuation of the Se sublattice across the interface.86 This mismatch is much smaller than the ∼8% mismatch for . Between the islands, Mn incorporates into nanoridges, as indicated by the protrusions observed in empty-state STM and the rapid oxidation of Mn when laminar films are removed from the UHV growth chamber. Photoemission of laminar films shows that incorporation of Mn films shifts the bands upward, indicating a p-type surface consistent with substitutional rather than Mn occupying an interstitial intrinsic vacancy. Once the lattice strain is large enough, either due to increased lattice mismatch or to sufficient thickness at lower mismatch, the growth converts to a Stranski–Krastanov layer + islands mode. The islands are MnSe, while the terrace continues as (lightly Mn-doped) .
It is not known whether Mn-doped is magnetic, although two-dimensional grown on GaSe is known to be ferromagnetic at room temperature.89 Neither magnetic circular dichroism at , , nor variable temperature superconducting quantum interference device magnetometry exhibited any ferromagnetic signal above the noise level; however, since the measured films were only 1−2 nm thick and the experiments were ex situ (and hence the surface Mn oxidized), this does not rule out magnetism definitively.77 Nevertheless, the successful incorporation of low concentrations of Mn into is of interest for exploring potential p-type doping in intrinsic-vacancy selenide semiconductors, and the islanding at high concentrations highlights the importance of lattice mismatch in controlling disproportionation during strained heteroepitaxial growth.
2. Cr-doped Ga2Se3
In contrast to manganese, it is possible to fabricate thick (at least 20 nm), laminar films of chromium-doped on Si(001):As up to about 8 at. % Cr, after which Cr-rich islands nucleate.90 These laminar films are ferromagnetic at room temperature with a moment of about /Cr atom.16 Chromium doping has also been reported to result in room temperature ferromagnetism for a number of wide bandgap III–V (GaN, AlN)91 and II–VI (ZnTe)92 materials, where cation substitution leads to a tetrahedral local environment. Stoichiometric bulk compounds, however, place Cr in an octahedral environment while Ga is tetrahedral.93,94 The compound is antiferromagnetic ,93 while the compound is paramagnetic with some indication of ferromagnetic interactions within clusters.94
Scanning tunneling microscopy images of Cr: films deposited on Si(001):As at are shown in Fig. 19 for nominal Cr concentrations (a), 6% (b), and 9% (c) in the incident flux, where as for Mn.16,90 Photoemission measurements on flat films indicate that is equivalent to in .85 Low Cr concentration films [Figs. 19(a) and 19(b)] appear similar to pure (Fig. 15), except for bright spots attributed to empty Cr electronic states and a shorter length for a typical nanoridge. By [Fig. 19(c)], however, the nanoridge structure is replaced by a flat terrace containing widely spaced, tall islands that are surrounded by trenches which can reach to the underlying substrate. The terrace structure is similar to that of the initial As-rich (vacancy-poor) structure observed for very thin [Fig. 15(h)].
Islands like those in Fig. 19(c) are identified as a high Cr concentration Cr–Ga–Se compound surrounded by Kirkendall voids.90 Scanning Auger microscopy shows the islands to be Cr-rich, while photoemission spectroscopy shows them to be metallic; the facet orientation and alignment are consistent with the monoclinic or rhombohedral phases. Changes in the island structure with either an underlying buffer layer or an added capping layer of pure indicate extensive vertical interdiffusion of Cr and a saturation concentration in the terrace regions .90
Cr-doped films are ferromagnetic with a Curie point above room temperature [Figs. 19(d)–19(f)], with clear ferromagnetic hysteresis after subtracting the diamagnetic substrate [Fig. 19(e) inset]. The hysteresis loop, exhibiting a coercive field of about 200 Oe, is very similar at temperatures of 5 and 300 K [Fig. 19(e)]. The room temperature magnetic moment per Cr atom for this 3 nm-thick laminar film with is ; an additional paramagnetic component, possibly associated with As interdiffusion near the interface, is observed below 10 K [Figs. 19(d) and 19(f)]. The maximum observed moment per Cr was observed for , the saturation level for Cr in the terrace structure of Fig. 19(c); films with , where a majority of the Cr is in islands, exhibit a moment per Cr atom similar to that for laminar films with , indicating the islands are also magnetic.16 A difference in the moment during warming between field-cooled and zero-field-cooled laminar films persists up to the instrument limit [Fig. 19(f)], indicating a Curie temperature .
The local atomic and electronic structure of Cr in is deduced from absorption and emission spectroscopy (Fig. 20).16 The absence of a pre-edge x-ray absorption peak in Cr K-edge XANES [Fig. 20(a)], e.g., as in , shows that the Cr does not simply replace Ga in a tetrahedral site, as it does in GaN;95 rather, Cr is in a centrosymmetric environment. The XANES line shape and edge position indicate a local charge state similar to CrSe, with an additional component attributed to post-growth oxidation in this air-exposed, film [Fig. 20(a), bottom]. Chromium L-edge XANES [Fig. 20(c)] shows the same multiplet structure for Cr:as in , but with different relative intensities of the multiplet components. In situ Cr photoemission from laminar films without any oxide present also indicates , with an additional component observed in islanded films.16
Additional information is obtained from EXAFS [Fig. 20(b)].16 The predominant component in Cr K-edge EXAFS is for Cr in a locally octahedral environment with six Se neighbors at a distance , as modeled by either the bulk crystal (model A) or the defect structure proposed in Fig. 21 (model B). About 20% of the Cr is in a surface oxide with a bond length consistent with antiferromagnetic and not ferromagnetic ; 20% + 80% in a local CrSe environment is also consistent with the K-edge XANES line shape [Fig. 20(a)]. Ga K-edge EXAFS16,85 finds a Ga–Se bond length equal to that found87 for either pure or Mn-doped .
To assess the impact of increasing Cr concentration on the electronic states and local bonding, photoemission spectroscopy was obtained16 from a laminar film where the Cr concentration was varied with position in the range by repeated motion of a shutter across the Cr flux during growth [Figs. 20(d) and 20(e)]. The films are semiconducting and weakly p-type ; a new state attributed to Cr [shaded in Fig. 20(e)] arises below the Fermi level as the Cr concentration increases; this is the same energy range associated with Se lone-pairs.79 The bands for laminar films do not shift as Cr is added, indicating that Cr is substituting for the same-valence Ga, but islanded films are metallic, with the Fermi edge shifting 0.5 eV further from the valence band maximum, consistent with Cr in islands adding to, rather than just replacing, Ga.16 Increasing Cr concentration in laminar films also leads to a decrease in the Se component characteristic of Se atoms adjacent to two Ga and two vacancies (normally 1/3 of the total) and an increase in the component associated with an Se adjacent to a single vacancy, consistent with a decrease in both the number of vacancies and of nonbonded Se lone-pairs.
The deduced electronic and atomic structure of for is shown in Fig. 21.16 Figure 21(a) replots valence band photoemission [Fig. 20(e)] and x-ray absorption [Fig. 20(c)], aligned using Cr photoemission, reflecting both - and Cr-derived occupied states and Cr-localized empty states. The right side of Fig. 21(a) shows schematically the atomic origins of these states, assuming a crystal field that splits the Cr into triplet and doublet states. The octahedral environment places the states at lower energy, so the Fermi level lies between the and levels for the valence expected for substitution. Additional Cr incorporation beyond stoichiometric substitution would start to fill the states, making the system metallic and raising the Fermi level, as observed for higher-concentration islanded films. The empty states form the peak in x-ray absorption since the transition is symmetry forbidden in an octahedral environment. The strong overlap between and Se lone-pair states allows for indirect exchange between magnetically polarized Cr atoms in this dilute magnetic semiconductor. It is unclear whether the coupling is via itinerant holes or simply static overlap and hybridization; both mechanisms are feasible, as the average spacing between Cr atoms for is less than 1 nm.
The intrinsic vacancy structure of enables a straightforward conversion of a simple substitution on a tetrahedral site [Fig. 20(b) top] to an octahedral local environment for Cr [Fig. 20(b) bottom]. The Cr (dark gray, blue online) can shift (arrow) to the adjacent octahedral hole of the Se sublattice, which is also a tetrahedral hole of the cation sublattice. Two of the surrounding cation sites are already empty due to the alignment of vacancies in the epitaxial film, a third becomes empty when the moves, and the fourth can become empty by having the Ga atom diffuse to a nearby vacancy site (arrow). The EXAFS-derived bond length is about midway between the ideal 2.70 Å distance for the unperturbed octahedral hole in and the measured for pure, Cr-doped, or Mn-doped films, indicating local distortions of Se from the ideal fcc sublattice. The model in Fig. 20(b) is also consistent with the observed solubility limit of about one Cr per four intrinsic vacancies.16
Unlike Mn-doped gallium sesquiselenide, which leads to precipitation of the antiferromagnetic compound rock salt MnSe, chromium doping exploits the intrinsic vacancy structure of to form a novel magnetic structure. This octahedral local structure differs from that in either III–V or II–VI semiconductors with tetrahedral substitution, or other chalcogenides that form two-dimensional or topological magnetic materials (e.g., Refs. 96 and 97). It is also compatible with silicon technologies and holds promise for combining electronic and magnetic functions in a single device.
Over the past decade, the superlative properties of single-layer graphene have inspired aggressive exploration of other 2D materials, including transition metal dichalcogenides, but, more recently, broadening the larger family of layered materials as well as topological insulators to explore new physics. Nonetheless, silicon has been the backbone of modern electronics with its superior electronic, mechanical, and thermal advantages; therefore, the combination of Si with 2D materials represents an innovative approach to develop further miniaturized and multifunctional devices by harnessing the complementary advantages of both materials. Our research on III-Se materials over the past 25 years has mapped a plethora of heteroepitaxial structures obtained from III–VI heteroepitaxy on silicon. III–VI semiconductors contain intrinsic vacancies, flexible bonding constraints, and multiple sites for dopant incorporation. Bulk III–VI materials exhibit a wide variety of bonding environments, including both III–III and III–VI covalent bonds, and both chains and planes of weakly interacting anion lone-pair orbitals. These factors lead to a wide variety of stable local geometries, opening possibilities for new physics in novel nanostructures whose structure and properties can be controlled through hetero-epitaxial thermodynamics and kinetics. We have demonstrated the epitaxy of GaSe, Ga2Se3 and Al2Se3 on Si(111) and Si(001) and discovered novel means to passivate these Si surfaces, to eliminate interface reactions in subsequent growth of TiO2, to generate room temperature ferromagnetism in a nanoscale film, and to self-assemble low-dimensional nanostructures by exploiting their unique crystal structures. In addition, the close lattice matching of the prototype Ga2Se3 to silicon enables incorporation of the resultant optical, magnetic, or electronic nanostructures into silicon-based devices.
Many of the subjects described in this Review are relevant to more recent development of low-dimensional materials research and may help accelerate progress in those materials for electronics, photonics, sensing, and related applications in the coming years. These interdisciplinary fields are expected to experience continued growth and offer ample research opportunities for the foreseeable future.
We would like to thank the students and postdocs for their wonderful contributions to our research on selenides over the past 25 years. They are Lee E. Rumaner, Jennifer R. Grey, Uwe Hessinger, Aaron Bostwick, Scott A. Chegwidden, Ngigi wa (Isaiah) Gatuna, Shuang Meng, Jonathan A. Adams, Zurong Dai, Andreas Klust, Brett R. Schroeder, Diedrich A. Schmidt, Patrick Shamberger, Chih-Yuan (Claire) Lu, Taisuke Ohta, Eswar Venkatasubramanian, Esmeralda N. Yitamben, Xiaohao (Sam) Zheng, and Tracy C. Lovejoy. We also thank our University of Washington faculty collaborators: Scott Dunham, Alec Pakhomov, Qiuming Yu, and the late Samuel Fain. We are grateful to Keiji Ueno of Saitama University, Eli Rotenberg of the Advanced Light Source, Steve Heald of the Advanced Photon Source, and Wolfram Jaegermann and Andreas Klein of Technische Universität Darmstadt for their insightful advice and collaborations. This work was supported by a series of grants from the National Science Foundation (NSF) (Grant Nos. DMR-9801302, 0102427, 0605601, and 071641) as well as IGERT training grant (No. DGE 0504573). The authors also acknowledge generous gifts from the Micron Foundation and the International Fund from New Energy Development Organization (NEDO), Japan.
The data that support the findings of this study are available from the corresponding author upon reasonable request.
Marjorie A. Olmstead, University of Washington, received the AVS Peter Mark Memorial Award in 1994 “for elucidating the nature of semiconductor surfaces and the heteroepitaxial growth of insulating materials on these surfaces.” Marjorie Olmstead is a Professor of Physics and Adjunct Professor of Chemistry at the University of Washington, Seattle. Before joining the UW faculty in 1991, she was an Assistant Professor of Physics at the University of California, Berkeley, and a Faculty Scientist at the Lawrence Berkeley National Laboratory. Between receiving her Ph.D. from UC Berkeley Physics (1985) and joining the faculty there, she was a Member of the Research Staff at the Xerox Palo Alto Research Center. Her undergraduate degree in physics is from Swarthmore College (1979). She served as director of the University of Washington Dual-Titled Ph.D. Program in Nanotechnology from 2004 to 2017, and as principal investigator of its associated Interdisciplinary Graduate Education and Research Traineeship, “Building Leadership for the Nanotechnology Workforce of Tomorrow.” She currently serves as the Undergraduate Faculty Advisor and Associate Chair of Physics for the largest undergraduate physics program in the country, with over 210 physics bachelor’s degrees awarded in the 2019–20 academic year. Olmstead’s research has centered on the formation of interfaces between dissimilar materials and the intrinsic properties of the resultant nanostructures, with particular interest in materials that exhibit intrinsic vacancies in their crystal structure. Her current interests focus on interface control of conductivity in transparent oxides and on growth of semiconducting chalcogenide films. She and her co-author Fumio Ohuchi have been collaborating since 1993. In addition to the Peter Mark Memorial Award, Olmstead received an NSF Presidential Young Investigator Award (1987), the Maria Goeppert Mayer Award of the American Physical Society (1996), and an Alexander von Humboldt Research Award (2000). She is a fellow of both the APS and AVS, and has received the UW Society of Physics Students Undergraduate Teaching Award three times.